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By Morris Cohen, William P. Kimball

This paper deals with the calculation of the composition of the infinitesimal trace of alloy that transfers from the liquid to the solid state at each temperature during the equilibrium solidification of a binary solid solution, and which thereby enables the compositions of the coexisting liquid and solid phases to follow along the liquidus and solidus lines respectively during the cooling. A general expression for the composition (6) of the liquid-to-solid transfer at each temperature has been derived: d= Zlxl(xl - P) + Zsxs(P - xs)/Zl(xl - P) + Zs(P - xs) [1] where ZL and Zs are the slopes of the liquidus and solidus lines at the given temperature, Xl and 28 are the compositions of the stable liquid and solid phases at the same temperature, and P is the composition of the alloy under consideration. The 6 equation is valid for the process of equilibrium heating as well as cooling, and may be applied to vapor-to-liquid, vapor-to-solid, liquid-to-liquid, and solid- to-solid transfers if the corresponding equilibrium diagrams are known. The d-versus-temperature curves for two alloys P1 and P2 are shown in Fig. I. In each case, the liquid-to-solid transfer has the same composition as the stable solid phase only at the temperature of the beginning of solidification, and has the same composition as the stable liquid phase only at the temperature of the completion of solidification. At each intermediate temperature, the composition of the liquid-to-solid transfer differs by a finite amount from those of the two coexisting phases. Therefore, during equilibrium cooling, the transfers must merge into the solid phase from the liquid without forming a transition phase; otherwise the condition of equilibrium would be violated. At any given temperature during solidification, the composition of the liquid-to-solid

Jan 1, 1941

By Morris Cohen, William P. Kimball

This paper deals with the calculation of the composition of the infinitesimal trace of alloy that transfers from the liquid to the solid state at each temperature during the equilibrium solidification of a binary solid solution, and which thereby enables the compositions of the coexisting liquid and solid phases to follow along the liquidus and solidus lines respectively during the cooling. A general expressioll for the composition (6) of the liquid-to-solid transfer at each temperature has been derived: , = Zlxl{xl - P) + ZsXsiP - xs) [I] Zl(xl -P)+ Zs{P - xs) where ZL and Zs are the slopes of the liquidus and solidus lines at the given temperature, xl and xs are the compositions of the stable liquid and solid phases at the same temperature, and P is the composition of the alloy under consideration. The d equation is valid for the process of equilibrium heating as well as cooling, and may be applied to vapor-to-liquid, vapor-to-solid, liquid-to-liquid, and solid-to-solid transfers if the corresponding equilibrium diagrams are known. The d-versus-temperature curves for two alloys P1 and P2 are shown in Fig. I. In each case, the liquid-to-solid transfer has the same composition as the stable solid phase only at the temperature of the beginning of solidification, and has the same composition as the stable liquid phase only at the temperature of the completion of solidification. At each intermediate temperature, the composition of the liquid-to-solid transfer differs by a finite amount from those of the two coexisting phases. Therefore, during equilibrium cooling. the transfers must merge into the solid phase from the liquid without forming a transition phase; otherwise the condition of equilibrium would be violated. At any given temperature during solidification, the composition of the liquid-to-solid

Jan 1, 1941

By William P. Kimball, Morris Cohen

This paper deals with the calculation of the composition of the infinitesimal trace of alloy that transfers from the liquid to the solid state at each temperature during the equilibrium solidification of a binary solid solution, and which thereby enables the compositions of the coexisting liquid and solid phases to follow along the liquidus and solidus lines respectively during the cooling. A general expression for the composition (6) of the liquid-to-solid transfer at each temperature has been derived: d= Zlxl(xl - P) + Zsxs(P - xs)/Zl(xl - P) + Zs(P - xs) [1] where ZL and Zs are the slopes of the liquidus and solidus lines at the given temperature, Xl and 28 are the compositions of the stable liquid and solid phases at the same temperature, and P is the composition of the alloy under consideration. The 6 equation is valid for the process of equilibrium heating as well as cooling, and may be applied to vapor-to-liquid, vapor-to-solid, liquid-to-liquid, and solid- to-solid transfers if the corresponding equilibrium diagrams are known. The d-versus-temperature curves for two alloys P1 and P2 are shown in Fig. I. In each case, the liquid-to-solid transfer has the same composition as the stable solid phase only at the temperature of the beginning of solidification, and has the same composition as the stable liquid phase only at the temperature of the completion of solidification. At each intermediate temperature, the composition of the liquid-to-solid transfer differs by a finite amount from those of the two coexisting phases. Therefore, during equilibrium cooling, the transfers must merge into the solid phase from the liquid without forming a transition phase; otherwise the condition of equilibrium would be violated. At any given temperature during solidification, the composition of the liquid-to-solid

Jan 1, 1941

By William P. Kimball, Morris Cohen

This paper deals with the calculation of the composition of the infinitesimal trace of alloy that transfers from the liquid to the solid state at each temperature during the equilibrium solidification of a binary solid solution, and which thereby enables the compositions of the coexisting liquid and solid phases to follow along the liquidus and solidus lines respectively during the cooling. A general expressioll for the composition (6) of the liquid-to-solid transfer at each temperature has been derived: , = Zlxl{xl - P) + ZsXsiP - xs) [I] Zl(xl -P)+ Zs{P - xs) where ZL and Zs are the slopes of the liquidus and solidus lines at the given temperature, xl and xs are the compositions of the stable liquid and solid phases at the same temperature, and P is the composition of the alloy under consideration. The d equation is valid for the process of equilibrium heating as well as cooling, and may be applied to vapor-to-liquid, vapor-to-solid, liquid-to-liquid, and solid-to-solid transfers if the corresponding equilibrium diagrams are known. The d-versus-temperature curves for two alloys P1 and P2 are shown in Fig. I. In each case, the liquid-to-solid transfer has the same composition as the stable solid phase only at the temperature of the beginning of solidification, and has the same composition as the stable liquid phase only at the temperature of the completion of solidification. At each intermediate temperature, the composition of the liquid-to-solid transfer differs by a finite amount from those of the two coexisting phases. Therefore, during equilibrium cooling. the transfers must merge into the solid phase from the liquid without forming a transition phase; otherwise the condition of equilibrium would be violated. At any given temperature during solidification, the composition of the liquid-to-solid

Jan 1, 1941

By J. H. Henderson, L. A. Wilson, R. L. Gergins, R. J. Wygal, D. W. Reed

This paper presents a method of predicting the production history of an underground combustion recovery process. A rigorous solution of the thermodynamics and hydrodynamics involved is beyond the scope of this report. However, a practical scheme to appraise combustion recovery performance has been Worked out. By prudent assumptions, regarding the attainment of steady-state conditions, a trial-and-error solution, which satisfies both material balance and three-phase relative permeability requirements has been evolved and tested. As a combustion zone moves through a formation the interstitial water, the water of combustion and a portion of the oil wil1 be vaporized. These vapors will move downstream to a cooler region of the formation where they will condense. Part of the remaining oil will be displaced by the imposed gas drive and what remains behind will be consumed as fuel. The fluids will distribute themselves downstrean? in a manner which will satisfy the three-phase relative permeability characteristics of the formation. Three distinct zones will exist ahead of the combustion zone: (I) a zone termed a water bank which contains three mobile phases, oil, water and gas; (2) a region containing two mobile phases, oil and gas, called the oil bank; and (3) a zone having the fluid saturution distribution which existed prior to combustion. The mobile water will impose a water flood on the oil in the region containing three flowing phases. The process can be visualized as a simultaneous water .flood and combustion-supported gas drive. It is possible to estimate the saturation distributions in the water bank and oil bank as well as the production history for any combustion temperuture if the porosity, three-phase relative permeability characteristics, oil viscosity and initial saturations are known. This has been done for injection pressures ranging from I to 100 atm, oil viscosities of 10 to 1,000 cp and porosities from 20 to 40 per cent. Several of these calculations have been checked in the laboratory, with the result indicating rather good agreement between the predicted and observed production histories. The analytical and experimental techniques are described and the comparison of predicted and observed performance presented. INTRODUCTION Considerable research by the oil industry during the past 10 years has been directed toward the thermal recovery of crude oil. Of all schemes considered, the one receiving the most attention is in situ combustion. This is a process in which part of the oil is actually burned within the reservoir rock. The uniqueness of the process lies in its potential. Theoretically, this potential generates from the following: (1) it may open the door to vast reserves previously not economically producible, (2) it may induce increased production rates where low rates are now realized and (3) it may give high over-all recovery efficiency. The general idea of the process has received considerable attention in the literature '-' and the following brief description is intended for orientation only. In situ combustion can be visualized by considering a linear system, one end of which represents the injection well and the other the producing well. The formation at the inlet end is raised to ignition temperature by placing a gas or electric heater at the face of the formation or by the ignition of thermite, charcoal or similar material. While heating, the water and part of the oil in the heated region is vaporized and moved out into the reservoir by the injection of gas containing oxygen. In addition, part of the oil is displaced as liquid by the imposed gas drive. The oil not displaced is modified to an immobile, coke-like material (hereafter referred to as coke) by the high temperatures existing immediately ahead of the combustion zone. This coke is used as fuel, and the combustion front advances through the formation at a temperature which is probably in excess of 800°F. As the combustion front advancek it will drive ahead all interstitial oil not rendered immobile and all interstitial water plus the water formed by the combustion of hydrogen. As a result the sand behind the combustion zone is freed of all oil and water. Although the general description of the process has been outlined in the literature, there has been no consideration of the behavior of the fluids moving downstream from the combustion front. It is the purpose of this paper to present a technique

By V. B. Kurfman

Grain refinement of Mg-Al melts by carbonaceous additions has been attributed to nucleation by aluminum carbide. The effects of process and alloy variables are interpreted and predicted in terms of the dispersion and chemistry of this phase. The grain coarsening action of Be, Zr, Ti, R.E., chlorination, temperature extremes, and prolonged holding times is described. Measures necessary to insure an adequate dispersion of the catalyst are discussed. CARBON inoculation treatments have become fairly well known and used for grain refinement of magnesium alloys containing Al. Although there is general agreement that a nucleation process occurs, the process is not understood and the inoculants are used in a rather empirical fashion. The treatment is applied to the class of alloys containing 3 to 10 pct Al, i.e., AZ31A to AM100A. Typical methods involve melting, alloying, and adjusting the temperature to 1400° to 1450°F. Then 0.01 to 0.5 pct C as CaC2, C6C16, or lampblack is added by any convenient means, and the melt poured within 10 to 30 min. Investigators generally have been impressed by an assumed similarity of this refinement process to superheat grain refinement, which depends on heating approximately the same alloys to a temperature in the range of 1550" to 1650°F, then pouring promptly after the melt is cooled to the pouring temperature. Various predictions have been made that carbon refinement would replace superheating in commercial practice due to reduced process costs, but this replacement has not fully taken place because of production difficulties and conflicting observations. Davis, Eastwood, and DeHaven1 agree with Nelson2 and wood3 in suggesting that an excess of inoculant may be harmful. Wood however says that overtreat-ment is not a problem in production use of hexa-chlorobenzene inoculation, and Hultgren and Mitchell4 claim no evidence of harm from excess additions. Various grain coarsening reactions are known to occur, including the possibility of overtreatment mentioned above. Trace amounts of Be,2 Zr, and Ti may prevent refinement by either a carbon treatment or a superheat. Occasionally treatment with cl25 may cause coarsening, although the Battelle refinement process' uses a CC14-C12 blend. Grain coarsening also tends to occur on holding at temperatures below 1350°to 1400°F, especially after a superheat treatment, and for this reason Nelson2 stresses the desirability of a refinement method useful at lower temperatures for open pot melting practice. Since a carbon treatment can be made to work at temperatures below 1400°F, it seems desirable to investigate the mechanism of the refinement and the mechanisms of the coarsening reactions in order to establish control conditions for use in commercial production. The identity of the nucleating phase must first be established and then the factors affecting its chemistry and physical dispersion must be determined. THE IDENTITY OF THE NUCLEATING PHASE Davis, Eastwood, and DeHaven suggested that the nucleating phase in this system is Al4c3,1 but Mahoney, Tarr, and LeGrand8 disagree, largely because they found no evidence of the compound in alloys after carbon treatment and because there is no indication that aluminum carbide should be unstable over the temperature range used in the superheat treatment. This latter objection is based on the assumption that both the carbon treatment and the superheat treatment introduce the same nuclei. Electron diffraction studies have been made to identify the nucleating phase. Samples of grain refined A292 have been selectively etched SO that clean surfaces are obtained and so that secondary phases are in relief. Electron diffraction patterns from these surfaces have established that the carbon treatment of A292 introduces into the metal a large number of small, plate-like particles with a structure very similar to Al4C3. In most cases, the plate-like nature of the particles prevented positive identification but in the cases where the identification could be made the particles proved to be AIN A14C3. However, enough variation in lattice constants was observed so that all compositions from pure A14C3 to the 50:50 solid solution A1N.Al4C3 were probably present.14 In A14C3 and especially AlN.Al4C3 the A1 atoms occur in layers within which they have the same hexagonal symmetry and spacing as the Mg atoms in a single basal plane of a magnesium crystal. The solid solution spacing lies between the 3.16 of AIN and the 3.3? for Al4C3, in satisfactory agree-

Jan 1, 1962

By C. E. Ells, A. Sawatzky

The effect of neutron irradiation on tensile properties of the Zr-2.5 wt pet Nb-0.5 wt pet Cu alloy has been evaluated for integrated neutron fluxes up to 3 x 1020 n per sq cm (E > 1 Mev). Specimen temperature during irradiation was normally 300°C. The material was studied both in annealed and in quenched and aged conditions. When the alloy (fab-ricatedfrom sponge zirconium) is quenched from a temperature =40°C below the a + ß/ß transformation and aged 6 hr at 535oC, neutron irradiation up to the maximum studied has little effect 071 reduction in area, although there is a drop in uniform elongation and the yield strength is increased by = 25 pct. Conversely, irradiation of the alloy in an annealed condition can result in much greater changes in properties, with yield-strength increases of up to 200 pet. The dose dependence of irradiation hardening obeys the saturation equntion proposed by Makin and Minter to explain hardening in copper and nickel: Act = C[1 - exp(-Døt)/1/2 A rnodel is proposed to explain the effect of niobium content and metallurgical condition on irradiatiotz behavior. THE Zr-Nb alloys form a class of material susceptible to marked strengthening by quench and age heat treatment. With niobium concentrations as low as 2.5 wt pet, strengths can be developed which are double those of annealed zircaloy-2.1 Since the neutron-capture cross section of the Zr-2.5 wt pet Nb alloy is nearly identical to that of Zircaloy-2, significant gains in power-reactor neutron economy could be obtained by replacing stressed in-reactor components of Zircaloy-2 with the heat-treated Zr-Nb alloy. When the solution heat-treatment temperature is in the high (a + ß) phase, then the tensile properties of the quenched (and aged at 500°C) material have been shown to be relatively insensitive to neutron irradiation.1-3 Addition of small quantities of copper to the binary Zr-2.5 wt pet Nb alloy gives a useful reduction to the corrosion rate in air and carbon di- oxide,&apos; and a ternary alloy of composition Zr-2.5 wt pet Nb-0.5 wt pet Cu was chosen for one component in the Douglas Point Reactor. This application utilizes both the corrosion resistance to a moist air-carbon dioxide environment and the strength developed by a quench and age heat treatment. By suitable choice of solution heat treatment and aging temperatures, tensile properties of the ternary alloy can be made nearly identical to those considered as optimum for the binary alloy. One effect of the copper addition, however, is to give marked changes in the aging kinetics of the alloy.&apos; For this reason, it appeared that neutron irradiation might promote overaging in the ternary alloy, particularly at the elevated service temperature (=300°C). This paper describes the effect of neutron irradiation on the tensile behavior of the Zr-2.5 wt pet Nb-0.5 wt pet Cu alloy in several metallurgical conditions pertinent to its use as a reactor material. The metallurgical conditions included those for which aging occurs in the unirradiated material, and fully annealed conditions for which a higher degree of thermal stability could be expected. 1) EXPERIMENTAL Two batches of alloy were used, both made from sponge zirconium. The fabricator&apos;s analysis of principle constituents showed no significant difference between the two batches, Table I. The rod material was rolled to 1.5 in. diam, and then swaged to final sizes of 0.5 to 0.375 in. diam at nominal temperatures of 785" and 625°C for the AM and AS batches, respectively. Subsequent hydrogen analyses at CRNL confirmed that the hydrogen concentration of the as-received rod was indeed <20 ppm. However, it was found that the (a + ß)/ß transus for the AM material was approximately 25°C lower than that of the AS material, indicating rather more difference in oxygen concentration than given in Table I. Except for a small difference in yield-point behavior, the tensile prop-

Jan 1, 1965

By M. D. Banus, M. B. Bever, A. K. Jena

The heals of formation at 78", 195, and 273°K of the ternary compound h3SbTe2 based on the elements and based on the binary compounds In Sb and [inTe have been measured. The heats of formation at these temperatlcres of the binary compounds In Sb and In Te based on the elements have also been determined. Heal contents and free energies of the three compounds have been calculated from 0° lo 80I)°K. The free energies of joyrrzalion, heats of formations, and entropies of formation at 298°K have also been calculated. The results shown that the tertnary compound is metastable with vespecl to InSh and ln Te below 696 °K. bul is slable above that temperature. The weaker bonding of results in a positice entropy of formation which with incrensirzg temperature makes increasing negative conlvihtclions to the free energy and above 696°K renders the compound slable. THE ternary compound In3SbTez occurring in the quasi-binary system In Sb- In Te&apos; forms on cooling at 829°K by a peritectic reaction.&apos; Observations at 673" and 573 K have shown that this ternary compound decomposes slowly into the binary compounds InSb and1n~e.l&apos;&apos; It has not been possible to analyze the metastable behavior of the ternary compound because up to the present time data on its thermodynamic properties have been lacking. Some information on the binary compounds, however, is available. The heat of formation of InTe at 273°K and its free energy at 673°K are kn~wn.~&apos;~ The heats of formation of InSb at 78", 273&apos;, 298", and 723°K have been measured5-&apos; and its heat capacity between approximately 12" and 300"Kg9l0 is also known. In the investigation reported here the heats of formation at 78% 195% and 273°K of the ternary compound In3SbTez have been measured. The heats of formation of the binary compounds InSb and InTe at these temperatures have been obtained by combining new calorimetric results with previously published data. The heat contents and free energies of the three compounds at various temperatures from 0" to 800°K have been calculated. Against the background of this information, the metastability of the ternary compound will be discussed. 1) EXPERIMENTAL 1.l) Preparation of Specimens. The materials used consisted of the elements indium, antimony, and tellurium, the binary compounds InSb and InTe, and the ternary compound In3SbTez. The elements, obtained from American Smelting and Refining Co., had a nominal purity of 99.999+ pct. The compound InSb was Cominco semiconductor grade; the compound InTe was prepared from the elements by melting under a vacuum of 10-h m Hg followed by slow cooling. Three batches of specimens of the compound In3SbTez were used. One batch was prepared by melting appropriate amounts of the elements in an evacuated and sealed Vycor tube. The melt was held at approximately 100°K above the liquidus for about 8 hr, shaken repeatedly, and quenched into a mixture of ice and water. The specimen was annealed in vacuum at 760°K for 4 weeks. In preparing a second batch, a mixture of the component elements was melted and quenched in water. The resulting ingot was powdered. The powder was pressed into pellets, which were annealed in vacuum at 748" to 773°K for 4 weeks. A third batch was prepared in the same manner as the second, except that the starting materials were InSb and InTe rather than the elements. Metallographic examination of samples of the three batches and X-ray diffraction examination of a sample of the second batch did not reveal evidence of microsegregation or a second phase. The results obtained with the three batches showed no systematic differences. 1.2) Calorimetry. The calorimetric method has been described in detail." Samples of the compound In3SbTez, mechanical mixtures of InSb and InTe in the proportion of 1:2, and mechanical mixtures of indium, antimony, and tellurium in the proportion of 3:1:2 were added to a bismuth-rich bath at 623°K in a metal-solution calorimeter. These additions were made from 78°K (liquid nitrogen), 195°K (dry ice and acetone), and 273°K (ice and water). The heat effects of the additions were measured. The difference in the heat effects of the additions of a compound and the additions of the mixtures of its constituents, adjusted for differences in the concentration of the bath, is the heat of formation of the compound. In the concentration range not exceeding 1.5 at. pct solute, the heat effect of the additions was a linear function of concentration. The heat of formation refers to the temperature from which the additions were made, namely, 78", 195", or 273°K. Several calibrating additions were made in each calorimetric run. The calculated heat capacity of the calorimeter and hence the calculated heat effects of the additions of samples depend upon the values adopted for the heat contents of the calibrating substance. In this investigation bismuth at 273°K was used and a value of 4.96 kcal per g-atom was taken for (HGZ3"k . 2) RESULTS AND DISCUSSION 2.l) Heats of Formation. The heats of formation of the ternary compound In~SbTez from the component elements indium, antimony, and tellurium and from

Jan 1, 1968

By M. E. Fine, A. A. Hendrickson

The critical resolved shear stress and the strain rate dependence of the .flow stress are reported for Ag base A1 single crystals up to 6 at. pct A1 over a temperature ralzge of 4.1° to 470°K. At room temperature the solid solution strengthening is attributed brinciballv to an increase in the dislocation density on alloying. At 42°K the alloys are relatively stronger because the dislocations widen on alloying and are more difficult to cut. In a previous paper, the authors reported the characteristics of strain aging in Ag base A1 alloys.&apos; The variation of the magnitude of the strain-aging yield drops on aging time, aluminum content, and testing temperature strongly indicated that the phenomenon is due to the Suzuki mechanism. In the present paper, the critical resolved shear stress (CRSS) and the strain rate dependence of the flow stress are reported for Ag base A1 single crystals as functions of testing temperature (4.2" to 470") and A1 content (0 to 6 at. pct). Observations on the substructure and measurements of the effect of crystal growth rate are also given. The experimental re- sults appear explainable considering that aluminum additions increase the dislocation density and dislocation width. EXPERIMENTAL METHOD The methods of specimen preparation and tensile testing were described previously.&apos; Briefly, single crystals were grown from the melt in a vacuum (1 x lom5 mm Hg), shaped into tensile samples by a combination of abrading and acid etching, annealed at 850°C in a vacuum (1 x 10" mm &Z) for several days, and furnace cooled to 200&apos;. The testing was done on a model TM Instron using friction grips and a strain rate of about 5 x 106 per sec unless otherwise noted. The various testing temperatures were obtained by immersing the specimens in the following media: 1) 4.2&apos;: liquid helium using the apparatus of G. Byrne;&apos; 2) 77°K: liquid nitrogen; 3) 200"K: dry ice and acetone; 4) 415&apos; K: ethylene glycol; 5) 460&apos; \ to 480&apos; : molten Stanalax. In the determination of the CRSS above 296&apos; K, a ratio method was used; the crystals were strained initially at the higher temperature, 0.1 pct shear strain, and then tested at 296&apos; . The ratio of the flow stress between the two temperatures was then multiplied by the average CRSS at 296&apos; K to obtain the value of the CRSS at the high temperature. (The test at 296&apos; showed a yield point resulting from strain aging; the flow stress at 296&apos; K was taken as the lower yield stress.) The ratio method, we believe, is more accurate than using different crystals since

Jan 1, 1962

By T. H. Alden

BaCKOFEN, Avery, and Turner have shown"2 that the large tensile elongation in superplastic metals is correlated with a high strain-rate sensitivity of the flow stress. At present, the reported superplastic materials are complex eutectic or eutectoid systems with a large solubility range in at least one of the terminal phases. This note describes the behavior of a solid-solution alloy (at 22°C)5,= Sn-1 wt pct Bi, which has been processed in order to produce a fine grain size.2,3,6 The chill-cast, 1/2-in.-diam ingot was homogenized for 1 week at 130°C, extruded at O°C to 0.083 in. diam, quenched within 15 sec, and stored at -196°C. Samples were warmed to room temperature 16 hr prior to mechanical testing or metallography. During tension testing at room temperature, it was found that the flow stress reached a steady-state value which depended on the imposed elongation rate and was independent of strain to an error of about ±5 pct. In Fig. 1 is plotted log a vs log e where a is the measured steady-state flow stress and i the imposed strain rate. These quantities were calculated using "instantaneous" area and length to an accuracy in the sample diameter of ±5 x 10-4 in. The log a-log 1 curve is sigmoidal. The strain-rate sensitivity, which may be given by d log old log 6,&apos; has a maximum value of about 0.48 between i = 5 x 10-4 and 5 x 10-3 per min. The value is well above frequently measured values for metals at creep temperatures (0.1 to 0.3) and comparable to the rate sensitivity for Pb-Sn solder: which has been shown to be superplastic.2,3 Note, however, that below or above the strain rate of 10-3 per min the rate sensitivity of Sn-1 pct Bi decreases and at high rates is more nearly like that of conventional metals. The grain size of this material is relatively uniform, Fig. 2, and may be specified by an average linear intercept i.&apos; Grain growth occurs at room temperature but after 16 hr, L = 5. Tensile elongation was measured on finer-grain-size specimens, which had been at room temperature 15 min prior to testing. After 15 min, the samples are fully recrystallized. Three strain rates, 5 x 10-3, 10-2, and 2.5 x 10-2 per min, were employed for which the measured elongations varied from 425 to 500 pct. These results would seem to justify the term superplastic, applied to this alloy, although the ductility is inferior to that of the Sn-Bi eutectic.3 The Sn-1 pct Bi alloy, which is a single-phase material with equiaxed, fine grains, is an excellent one on which to make an estimate of the Nabarro-Herring9 creep rate. According to the model of Avery and Backofen2 a substantial part of the deformation must be accomplished by N-H creep in the strain rate-temperature regime in which d log a/d log 1 = 0.5. The maximum N-H creep rate is given by9 L2kT where v is the atomic volume and D the self-diffusion coefficient. For this test v = 3.O x 10-23 cu cm, D = 1.4 e-23,300/RT = 2.5 x 10-l7 sq cm per sec,7 L = 5x 10+4 cm, T = 295°K, and a = 2000 lb per sq in. = 1.38 x 10&apos; dyne per sq cm. The calculated creep rate is E= 4 x 10-11 per sec which may be compared with the measured creep rate at this stress, Fig. 1, of 6.7 x 10-5 per sec. The N-H creep rate is seen to be negligible. Earlier consideration2 of this question was in error because incorrect diffusion coefficients were used. The assistance of R. R. Russell in metallography is acknowledged gratefully.

Jan 1, 1967

By G. A. Beske, P. Hilsch, J. Wulff

The effects of the growth precipitates on the su-perconduching properties of an Al- 15 al. pel Zu alloy have been studied using magneization, transition lem-perature, and residual resistivity measurements. Aging at 230C for 1/2 to 11 he produces an increase in Hfp, Ike field of first penetration, and tut increase in trapped flux. The upper critical field, Hc2 remains constant for such aging tunes. Aging at 200°C produces a decrease in Hfp and an increase in trapped flux. The transition temperature remains constant at Tc - 1.227° ± 0.02°K for aging at 220°C for tims of from 1/2 to 10 hr, and for aging at 200°C it remains constant at Tc. = 1.18° i 0.02°K for times from 1/2 hv to 26 hr. The observed behavior indicates that the superconducting matrix when aged at 220°C behaves tike a type U su-perconducior, but when aged at 200°C behares like a type I superconductor. The magnetization changes after aging can be attributed to the growth and clutnge in distribution of the precipitate in a matrix of fixed composition. IT has been realized for some time that the superconducting properties of alloys are highly structure-sensitive. This is evident in the work of Mould and Mapother1 who examined the properties of an aluminum alloy. Precipitation effects have also been briefly mentioned in other studies.2-4 Bonnin and coworkers5 have observed a sharp inc rease in trapped flux in an A1-Mg alloy when the residual resistivity reaches a value that indicates it is a type II superconductor. Recently Blanc, Goodman, and Nemoz5p have reported on precipitation effects in A1-Ag alloys. Detailed studies of precipitation in Pb-Sn and Pb-Cd alloys made by Livingston6 serve to show that the upper critical fields are enhanced by quenched-in solute. Furthermore. at various stages in the precipitation process? the upper critical field is determined largely by the remanent solute. The observed magnetic hysteresis and trapped flux in such alloys result from the interaction between precipitate and the flux filaments in the mixed state. The flux trapping appears to depend largely on the precipitate distribution and is greater if the superconductor is type II than if it is type I. To examine the effect of precipitation in another alloy system, particularly one in which the composition of the superconducting phase does not change significantly during the course of precipitation, the Al-Zn system seems interesting, especially in view of the transition temperature and resistivity studies of Chiou and seraphim.&apos; Furthermore, the recent detailed electron-microscope studies of an Al-Zn alloy by Richards and Garwood8 provide a guide for relating superconducting properties to changes in structure. For these reasons, magnetization, transition temperature, and residual resistivity measurements for varying times of heat treatment at two temperatures were made with a -15 at. pct Al-Zn alloy. The results are reported in this paper. I) EXPERIMENTAL The alloy used was prepared by Cominco Products Inc. from 99.9999 A1 and 99.999 Zn. It was chill-cast, swaged, and drawn to 0.125-in.-diam rod. The analyzed composition was 15.6 at. pct Zn (31.1 wt pct). A jeweler&apos;s saw was used to cut the 0.125-in.-diam rod into the 0.875-in. lengths needed for magnetization measurements. The ends of the cylinders were slightly beveled to remove the sharp edges. A portion of the original 0.125-in.-diam rod was drawn into 0.010-in.-diam wire and used for transition temperature measurements. Specimens were suspended from wires within a vertical furnace during heat treatment. All were homogenized at 500°C for a minimum of 3 days and subsequently solutionized for times in excess of 6 hr. The solutionizing temperature was held constant to +3°C. Individual specimens were homogenized, quenched, and aged several times and no effects indicating loss of zinc were noted. The suspended specimens were quenched by cutting the wires and allowing the specimens to fall into a silicone oil bath held at the aging temperature. They were then isothermally transforked by aging in the same bath used for quenching. After aging for the required time (approximately 10 min aging was necessary before the upper critical field reached a fixed value), they were briefly (1 to 2 sec) dipped into an acetone bath at room temperature to remove adherent oil and then immediately transferred to liquid nitrogen where they remained until inserted into the helium bath (usually within an hour). The temperature of the aging bath was held constant to +3°C for longer aging times. The cryostat used for the measurement of superconducting properties consisted of outer and inner dewars for liquid nitrogen and liquid helium, Stokes 6-in. booster diffusion pump and Stokes Model 212 mechanical pump, temperature controller to measure and control the temperature of the helium bath, and a superconducting solenoid with inserted search coils to provide the magnetic field and measure the magnetic moment of the specimens. A schematic diagram of the cryostat is shown in Fig. 1. An ultimate temperature of 0.85°K was reached with this apparatus. The temperature controller used is similar to the one described by Blake and chase.&apos; A 0.5 pct Allen-Bradley carbon resistor and a nichronle heater (200

Jan 1, 1968

By J. W. Cuthbertson, R. A. Cresswell

THE constitution of Cu-rich alloys with 1.5 to 13.5 pct Sn and 0.25 to 3.0 pct Be and the precipitation-hardening characteristics of alloys with 1.5 to 13.5 pct Sn and 0.25 to 1.0 pct Be have been examined. The hardness and tensile strength of the alloys examined increase markedly after solution treatment at 700°C followed by heat treatment at temperatures between 200" and 450°C. By a combination of cold work and heat treatment, hardness values similar to those exhibited by commercial Be-Cu alloys containing 2.25 pct Be can be obtained with ternary alloys containing 9 pct Sn and 0.75 pct Be and containing 10 pct Sn and 0.5 pct Be. Marked hardening effects occur with alloys containing even less beryllium. By heat treatment alone, a hardness value of 310 diamond pyramid hardness can be obtained from an alloy containing 10 pct Sn and 0.75 pct Be. Preliminary tensile tests have shown that an ultimate tensile strength of 110,000 psi with an elongation of 23 pct is obtainable by precipitation hardening an alloy with 8 pct Sn and 0.75 pct Be. The precipitation-hardening process has been followed microscopically for certain alloys and the inference is that, while the initial hardening effect is probably explained by the precipitation of the ß phase of the Cu-Be system, further hardening, proceeding at a much slower rate, also occurs, apparently as a result of precipitation of phases of the Cu-Sn system, particularly precipitation of the 6 phase at temperatures below 350". The presence of the e phase of the Cu-Sn system in certain alloys at temperatures below 350°C has been confirmed. Tin-bronzes are widely used in engineering applications where a combination of high strength and good resistance to corrosion is wanted. The maximum strength is induced in these alloys by cold working, and it would be an advantage for many purposes if high strength could be achieved alternatively by an age-hardening process. While Cu-Sn alloys have a good fatigue resistance they can be surpassed in this respect by Cu-Be, but the use of the latter alloy is limited by its high cost. If, by adding beryllium to tin-bronze, the properties of the respective binary alloys could to some extent be combined, a most attractive alloy should result. As pointed out by Raynor,¹ beryllium is on the borderline of the zone of favorable size factors for copper, and the solid solubility of beryllium in copper is consequently much more restricted than if the size factor were strongly favorable. The size factor is sufficiently favorable, however, to permit an increase in solid solubility with rise in temperature, and there is thus a composition range in which CU- Be alloys are susceptible to hardening by precipitation heat treatment. Although the a phase of the Cu-Sn system is similarly susceptible to precipitation treatment, the time necessary to establish equilibrium in commercial alloys of this type is usually so great that age hardening becomes impracticable. The addition of beryllium to Cu-Sn alloys would appear to offer a means of conferring on the latter useful age-hardening properties. Masing and Dahl² and others have, in fact, shown that the addition of beryllium to Cu-Sn a solid solutions renders these alloys susceptible to precipitation hardening and after such hardening confers on them an encouraging improvement in physical properties. If this improvement could be achieved by the addition of substantially smaller amounts of beryllium than are customarily found in binary Cu-Be alloys, the ternary alloys should possess economic advantages which might make them more attractive than the binary alloy for some applications. Binary Systems Copper-Tin: The constitution of these alloys is now reasonably well known and is summarized in the equilibrium diagram published by Raynor.³ The following observations, due to Raynor,¹ on the structure of those phases of the Cu-Sn system that are likely to be found in the ternary alloy system will facilitate the subsequent discussion on the examination of that system. The ß phase is an electron compound at the electron-atom ratio 3:2 and has a body-centered cubic crystal structure. This phase is stable only down to 586°C, at which temperature it decomposes eutectoidally into the a and y phases. The y phase has a structure that is also based on the cubic system. This phase is stable down to 520°C, at which temperature it decomposes eutectoidally into the a and d phases. The d phase is an electron compound (Cu³¹Sn8) which has a crystal structure analogous to that of 7 brass. This phase is stable from 590" to 350°C; on prolonged annealing at the latter temperature it breaks down into a mixture of the a and E phases. The e phase is an electron compound (Cu³Sn) having the electron-atom ratio 7:4. Its structure may be regarded as a superlattice based on the close-packed hexagonal system. This phase is stable from 676°C to room temperature. The primary solid solubility of tin in copper increases to a maximum of 15.8 pct as the temperature falls from that of the peritectic reaction to 586°C. The solid solubility remains constant from 586" to 520°C. At lower temperatures the solubility decreases progressively. Below 350°C the fall in solubility is pronounced and is associated with the precipitation of the e phase. This precipitation is very sluggish and does not normally occur under service conditions. Copper-Beryllium: The Cu-Be system has been investigated by Borchers&apos; and others. Raynor5 summarized the present state of information on it.

Jan 1, 1952

By V. J. Sikora, T. S. Hutchinson

This paper presents a method for making a water-rlrive ana1gsis without prior knowledge of aquifer geometry and uniformity using a standard desk calculator. Although it is necessary to know the initial oil in place to use this method, this is a minor limitation to the scope of (he method since in most reservoirs it is possible to obtain reasonably accurate volumetric oil-in-place estimates. An equation is developed which relates pressure at the water-oil contact to the water influx rate as a function of time by a factor called the resistance function. The resistance function introduces the composite effect of the aquifer geometry and flow resistance disrribution. It is pointed out that the characteristic shape of this function makes it possible to start with on approximation of the function and successively improve the approximation until the correct resistance function curve is obtained. In this fashion water-drive estimates can be made without the limitation of assuming simplifed aquifer shapes and flow distributions. This is the novel feature of this development. Methods are given for extrapolating the final curve to calculate future aquifer behavior. Equations are developed for adjusting the pressures and water influx rates where it appears that possible errors in these quantities make it difficult or impossible to obtain a useuble resistance function curve without this adjustment. Application of the pressure build-up analysis techiiique to estimate. some of the aquifer properties is also presented INTRODUCTION A great many oil pools are the result of oil accumulation in some type of trap in an otherwise large and continuous porous stratum. The void space of this stratum outside of the oil pool itself is filled with water or brine. In analyzing performance of the oil pool surrounded by this water aquifer, it is quite necessary. in most cases, to include behavior of the water. When the pressure at any point in a fluid system is lowered, such as by opening a well in an oil sand, fluids in the immediate neighborhood of this point will begin flowing towards this lower pressure sink. As pressure in this area drops due to flow towards the sink, fluids from farther out will start to flow towards the lower pressure. As more and more fluid is removed from the system, the distance from the sink or well within which flow is occurring will continually increase; that is, thc region of disturbance will grow. If some rigid boundary such as a fault is reached by the disturbance, this area will cease to grow; but if some movable boundary is reached, such as a water-oil contact, the area will grow on out into the water, although rate of growth may and almost always does change. The relation between amount of pressure drop and amount of fluid flowing at any point and at any time in an aquifer depends on such factors as compressibility and viscosity of the fluids, the porosity and permeability of the rock, geometry of the whole system and withdrawal rate or pressure drop. With these factors known, it is theoretically possible to calculate the pressure-flow behavior of the system, However, in practice true solutions to the problem are next to impossible due to complexity of reservoir systems. A number of approximate methods of solution have been developed based on various simplifying assumptions. One frequent assumption is that the water-oil contact can be located and equations defining oil reservoir and aquifer solved separately by assigning values to various parameters so as to match past pressure and production history. Usually the properties of the aquifer are not known, since few, if any, wells are drilled through the aquifer; but the water influx and the pressure at the reservoir boundary are known over some time. If it can now be assumed that the aquifer is circular, pie-shaped, or linear, and that it has uniform properties, it is possible to fit a theoretical dimensionless curve to the past aquifer performance history and therefore calculate future water drive. These theoretical curves are available in the literature,1,2,3 Of course, the assumption of uniform aquifer properties is almost always somewhat in error. Fortunately, however, moderate variations from the average have little effect on behavior of the system; and, hence, the best fit of a theoretical curve is frequently satisfactory. In other cases variations in aquifer properties and geometry are large enough that none of the available theoretical curves will give an acceptable fit of the data. In these cases, methods of fitting the data with various electrical analyzers have been developed.1,2&apos; It is the purpose of this paper to present an approach

By P. G. Shewmon, H. E. Collins

A study has been made to determine the sites at which sulfur adsorption occurs on copper surfaces. measurements were made of the relative torques, Ys, at the intersection of twin boundaries with surfaces near the three low-index orientations, i.e., (100), @lo), and 011), over a range of H2S/H2 ratios. HZS concerztvations j&apos;ro~n 3 to 1500pp)n between 830" and 1050°C were used. It is concluded that sulfur adsorption occurved preferentially though not exclusively at edge sites near the (100) and (110) surfaces in the HzS range — 700 Ppm giving rise to negative torques near these orientations. Beyond this HzS range, adsorption occurred at all sites. Near the (111) surface, 7/y little with HzS concentration up to approxiwzately 75pptn. Above this range, the results indicate adsorption is occurring OH both terrace and edge sites. SCIENTIFIC interest in surfaces and their interactions with a gaseous environment dates back to the beginning of the 19th century. The scientific luminaries of that period—Faraday, Maxwell, Rayleigh, Dewar, and Gibbs—were already concerned about such processes. However, it has only been within the past several decades that adsorption on metal surfaces has been actively studied. This increased interest in adsorption has been brought about by the advent of new and improved experimental techniques and apparatus, e.g., ultrahigh vacuum, and field-emission and ion microscopes. However, most of the work done using these techniques has been carried out at low temperatures. When adsorption studies have been made at either low or high temperatures, they usually gave no indication of the particular surface orientations or type of sites on which adsorption was occurring. In the last few years, there have been a series of studies in which the surface tension, y,, and/or its derivative with respect to orientation, 7, have been studied as a function of orientation and atmosphere.&apos;-7 Nearly all of the work on the relative torque,* ~/y, silver annealed in hydrogen and air.6 Recently Winterbottom and Gjostein" have used a modified and more accurate Mykurian method to determine the y plot of gold in hydrogen The only work in which T/~, has been measured over a range of chemical potentials for a given solute, p2, is that of Robertson and shewmon7 on the Cu-0 system. They measured T/Y, vs Po, (10"" to 10- l3 atm) at 1000°C in various mixtures of Hz0 and HZ. From this work they estimated the value of p2 at which one half of the surface sites are occupied with oxygen, pg, as being in the range 10- l6 to 10- l5 atm of oxygen. They also found that increasing Pa increased the magnitude of ~/y, near the (111) and (100) orientations. This indicates that oxygen is not adsorbed preferentially at step edges, but uniformly over all surface sites. In addition, they did one experiment on sulfur adsorption on copper surfaces, which indicated that sulfur adsorption decreases ~/y, near the (100) orientation, while not affecting ~/y, near the (111). This could be interpreted as indicating that sulfur adsorbs preferentially at step edges near the (100). In this paper the primary objective of the work has been to carry out a study of sulfur adsorption on copper surfaces over a range of temperatures and p,. In conjunction with this work, thermal grooving at grain boundaries has been examined as a method of determining the effect of sulfur adsorption on y,. METHODS Ideally, one would like to have information on the quantity of solute adsorbed on a surface and the types of sites at which it is absorbed as a function of p2. The total quantity adsorbed or the surface excess is given by the thermodynamic equation Thus data on the variation of y, with pz indicates the value of p2 at which adsorption becomes appreciable and the quantity adsorbed. The type of adsorption site is more difficult to deduce but information on this can be obtained from the variation of rz with 8, the angular deviation of the surface orientation. This is obtained from the thermodynamic equationlg Data on t and ys as functions of p2 have been obtained by the following methods. 1) Twin Boundary Grooving—By determining the effect of adsorption on the torque, 7, where T is the variation of surface energy, y,, with orientation, it is possible to obtain some indication as to the preferred sites of adsorption. Experimentally, the torque value measured is the relative torque, 7/ys The twin boundary grooving technique suggested by Mykura&apos;&apos; was used in this study to determine near the three low-index orientations— (loo), (110), and (111). Mykura&apos;s equation relates 7 /yS to measurements of the di-

Jan 1, 1967

By K. A. Jackson, G. A. Chadwick, A. Klugert

The diffusion field ahead of a lamellar interfnce is analyzed using an electrical analog. A self-consistent solution is obtained for the shape of the interfnce and the diffusion field by an iterative process. The solutions presented here are for a 50-50 eutectoid or eutectic, The shape of the interface is found to he independent of growth velocity and lamellar spacing, and to depend on the relative values of interfacial free energies at the phase houndaries . The mode of growth of lamellar eutectics and eutectoids has been a subject of much interest for many years.1-4 Mehl and Hagel 1 have shown photomicrographs taken by Tardif when he attempted to determine experimentally the shape of an advancing pearlite interface; the results are completely ambiguous. Brandt&apos; and schei13 have made approximate calculations of the composition ahead of a lamellar growth front. The shape of the advancing front and the composition distribution ahead of the front are difficult to calculate because one depends on the other. It is the purpose of the present paper to describe a method by which this calculation has been done. Lamellar-eutectic growth usually occurs under conditions where the growth is fairly rapid, and the interface temperature is close to the eutectic temperature. The growth rate is usually determined by heat flow. Eutectoid growth, on the other hand, can best be studied by quenching to some temperature, and allowing growth to proceed isother-mally. In both cases the growth is believed to be controlled by diffusion* rather than by the atomic kinetics of the transformation. This being the case, a single treatment of the diffusion equation will apply to both cases, provided the region of the interface in a eutectic may be considered to be isothermal. If a part of the interface could appreciably change its thermodynamic driving force by advancing ahead of or lagging behind the mean interface, then the two cases would not be similar. Eutectics normally grow in temperature gradients of the or-der of a few degrees per centimeter. The normal eutectic spacing is the order of a few microns. Part of the interface would have to extend many lamellar spacings ahead of the mean interface before it experienced sensibly different conditions. The interface temperature is usually a few tenths of a degree below the eutectic temperature so that temperature differences of the order of one-thousandths of a degree (a displacement of one lamellar spacing) would be unimportant. Protrusions large compared to the mean spacing do occur when one phase only grows into a eutectic liquid. This is usually a dendritic type of growth, and easily distinguishable from the lamellar mode of growth. A single treatment of lamellar growth will apply equally well to both eutectic and eutectoid decomposition. At the interface, which as shown above is essentially isothermal, the difference between the equilibrium eutectic temperature Teu and the actual interface temperature Ti, can be divided into two parts: 1) the composition varies across the interface, so that the local equilibrium temperature is not Teu; and 2) the interface is curved, so that the local equilibrium temperature is depressed according to the Gibbs-Thompson relationship. This undercooling can be written as Teu-Ti =?T = mAC(x) + a/r(x) [1] where ?C(X) is the departure of the composition at a point x on the interface from the eutectic composition, see Fig. 1, r(x) is the local radius of curvature at a point x on the interface, m is the slope of the liquidus line on the phase diagram, and a is a constant given by where s is the interfacial free energy, TE is the equilibrium temperature, and L is the latent heat of fusion. The calculations in this paper will be made only for the case where the phase diagram is symmetric, that is, the eutectic occurs at 50 pct, the liquidi have the same magnitude slope m at the eutectic temperature, and C,, the amount of B rejected when unit volume of a freezes, see Fig. 2(a), is the same for both phases. As shown in Fig. 2(b), the composition ahead of the a phase will be rich in B, the composition ahead of the ß phase will be rich in A. The composition at the phase boundary is the eutectic composition. The difference between the local liquidus temperature and the actual tempera-

Jan 1, 1964

By B. E. Terkelsen, B. J. Piearcey

A study has been made of the effect of unidirectional solidification on the creep behavior, stress-rupture properties, and thermal shock resistance of four nickel-base superalloys. The alloys Mar-MZOO, B-1900, IN 100, and TRW 1900 show improved rupture ductility and thermal shock resistance when tested with the columnar grains parallel to the major stress axis during test. The relative improvement in creep-rupture properties depends on the intrinsic strength of the alloy, a property which depends on composition, heat-treatnient, and crystallographic orientation. The data clarifies some of the factors affecting the properlies of the cast nickel-base superalloy. COMPONENTS designed for high-temperature use have, in recent years, been fabricated by precision casting techniques using nickel-base alloys developed specifically for use in the conventionally cast condition. This development was a result of both the increasing complexity of the component and the recognition that high-temperature strength was incompatible with workability. The use of castings can be economically favorable, but, more important, the recent complex designs of air-cooled gas turbine blades and vanes in alloys which possess the necessary high-temperature strength cannot be forged. Common modes of failure of high-temperature components are excessive creep, creep rupture, and thermal fatigue. If rupture occurs then the mechanism is usually by inter crystalline cracking along those grain boundaries oriented transverse to the major stress axis. In the stronger alloys, rapid propagation of intercrystalline cracks result in apparent premature failure demonstrated by the absence of third stage creep in a creep-rupture test and low rupture ductility. Tensile ductility shows a similar trend, that is, decreasing ductility with increase in strength. It is not surprising, therefore, that increases in creep strength have only been obtained with a loss in resistance to thermal shock, a property which shows a dependence on tensile ductility. VerSnyder and ~uard&apos; showed that the application of unidirectional solidification to a brittle Ni-Cr-A1 alloy both improved the rupture life of the alloy and increased rupture ductility. Since this casting method results in columnar grains, it appeared that the solution to the lack of ductility in nickel-base alloy components was the elimination of the source of failure, namely the transverse grain boundaries. This concept was recently developed2 to produce cast-to-size gas-turbine blades and vanes consisting entirely of columnar grains oriented parallel to the major stress axis of the component. Not only did the process impart increased rupture life and ductility to the components, but it also increased their resistance to thermal shock. During this development, the effect of unidirectional solidification on the properties of several alloys was investigated. The resulting data allows certain conclusions to be drawn regarding the factors affecting the creep and stress-rupture properties of the cast nickel-base superalloy and also its resistance to thermal shock. Detailed information was obtained on the alloys Mar-MZOO, B-1900, In 100, and TRW 1900. The composition of the alloys studied are shown in Table I. MATERIAL PREPARATION AND TESTING The conventionally cast alloys are cast as +-in.-diam bars in innoculated shell molds under conditions designed to maintain control of grain size. Unidirec-tionally solidified casting were produced in the form of 3-in.-diam ingots. Specimens from the latter were machined such that the axes of the columnar grains were parallel with the axis of the test bar, except in certain cases where the transverse properties of the material were being evaluated. The specimens were tested in an axial loading creep machine.&apos; Temperatures were measured with thermocouples, attached just outside each end of the gage length, to assure uniform temperature along the entire specimen. Chromel/alumel thermocouples were used up to 1800"F and Pt/Pt-10 pct Rh thermocouples above 1800"F. The specimen temperature was recorded and maintained within i2"F throughout the test. Creep extensions were measured automatically with extenso-meters attached to ridges on the specimen and recorded continuously using linear variable differential transformers and multipoint recording equipment. MICROSTRUCTURE The microstructures of transverse sections of the four alloys studied are shown in Figs. 1 and 2. Fig. 1 shows optical micrographs of the alloys, both in the conventionally cast and unidirectionally solidified conditions. Each alloy displays a cored dendritic structure, a distribution of MC carbide, and the presence of the y — y&apos; eutectic constituent.~ In the case of IN 100 and B-1900, Fig. 1 indicates that in these specimens the eutectic is degenerate. That is, it consists only of y&apos; pools indicating that the y constituent has been solutioned during the cooling period after solidification. The major effect of unidirectional solidification on the microstructures of the alloys is a tendency to promote a more dendritic form of MC carbide and an increase in dendrite arm spacing, both effects being a result of the slower solidification rate in the unidirectional solidification technique. V, lec tron micrographs of transverse sections of the four alloys in the two cast conditions are shown in

Jan 1, 1968

By Robert B. Sosman

ONE of the most neglected fields for physicochemical education as well as for research is that of high-temperature phenomena. Few universities or technical schools give instruction in the physical chemistry of substances at high temperatures. On the contrary such instruction is almost entirely confined to the region between the freezing and boiling points of water. This field should be extended to a much wider range of temperatures, both low and high. Such an extension is logical, sound, feasible, and practical.

Jan 1, 1943

By P. P. Gregory

ESTIMATION of re- serves in prorated sand fields has been discussed by S. A. Judson, H. D. Easton, Jr., and W. A. Schaeffer, Jr., in a paper that appears in Vol. 114 (1935), of the A.I.M.E. TRANSACTIONS. The purpose of this present paper is to extend the investigation to prorated flush limestone fields and to outline a method of procedure in determining the amount of oil originally in place and the amount of oil that can be recovered from such fields. Although it is possible to use the decline curves to estimate the amount of recoverable oil in certain limestone fields that are allowed to produce the oil at or near capacity production, such a method cannot be used in fields under severely restricted production rates. The problem is exceedingly complex due to fluctuation in the allowable rates of oil production over a wide range and at frequent intervals, and especially because of the heterogeneous nature of the reservoir itself. A well completed in a large solution cavity, capable of producing great quantities of oil per day, may be offset by a well of very limited capacity. Also, it is di5cult to arrive at an average figure for porosity in a given limestone reservoir due to abrupt variations in the character of limestone, in many instances even between two ad- joining wells. It is a safe statement that uniformity of porosity and permeability in limestone reservoirs .is nonexistent. It is evident, therefore, that the method of estimation of re- serves on the basis of per-acre content can give only approximate results.

Jan 1, 1935

By L. D. Lash, J. R. Ross

Discovery of scandium in the Vitro solvent extraction plant for uranium led to commercial recovery of the byproduct. Micro amounts of scandium were extracted with uranium by dodecyl phosphoric acid, but failed to follow ura-niuminto the hydrochloric acid strip solution and were eventually concentrated in the organic extractant. A fluoride strip system was developed to recover scandium from the solvent in concentrated form. High purity scandium oxide was then prepared in multi-pound lots by chemical separation techniques. The plant recovery operation, final product purification, and analytical procedures are described. Scandium is a pseudo-rare earth which is truly rare and expensive. It has special properties which may make it desirable even at the present price of $2750 per lb. Recently the price was lowered from $5000 per Ib which had prevailed since 1952, and it is anticipated that usage will be stimulated because of the lowered price. This metallic element is distributed widely in trace amounts in the rocks of the earth&apos;s crust. In addition to micro-amounts found in most uranium ores, concentrations of 50 to 100 ppm have been found in Colorado ferberite ore, lateritic nickel ore, heavy sandstone from Utah, Wyoming, and New Mexico, and various zircon sands and monazites from the western part of the U. S. Scandium has been found as an essential constituent of very few minerals, the most important being thortveitite, a scandium silicate. At the present time, this mineral commands a price more than twice that of gold. Chemically, scandium is in group 3A of the periodic table. It is closely associated with yttrium and the rare earths which it strongly resembles in its reactions. It has a valence of three and will form an insoluble hydroxide, fluoride, or oxalate under proper conditions, similar to the rare earths. Eight years before its discovery, scandium was predicted by Mendeleef who ascribed his eka-boron with properties nearly identical to the actual properties of scandium. Its position as the element with Atomic No. 21 and Atomic Weight 44.96 places it between calcium and titanium in Period 4 of the chart. The successful commercial application of scandium must depend on unique properties to justify use of this high cost substance. Greater consumption would allow additional price reductions, but scan- dium will remain a relatively high-priced commodity because of its inherent scarcity. Research reportedly has been oriented toward use of scandium as an agent to achieve high-temperature, low-density alloys with magnesium and tantalum.&apos; Also, the U.S. Air Force sponsored a project for production of research amounts of pure scandium metal.&apos; The oxide has been used experimentally in electronics, ceramics, and metallurgy. The various aspects of scandium are well covered in recent publications.s-5 CONCENTRATE RECOVERY Small amounts of scandium exist in uranium ores and are dissolved during an acid leach yielding up to 0.001 gpl Sc2O9. When Vitro converted their uranium plant from phosphate precipitation to solvent extraction: scandium was found to follow the uranium into the dodecyl phosphoric acid solvent (DDPA). However, scandium did not strip with uranium from the DDPA in hydrochloric acid, but remained in the solvent. Therefore, a concentration of scandium built up in the organic phase. Residues from the solvent were spectrographed by the U.S. Bureau of Mines and found to contain scandium. As a result of this discovery, provisions were made in the plant for recovery of scandium-bearing concentrates. Solvent extraction of uranium was added to a conventional acid leach process in a typical installation? The ore was crushed and ground, and then leached with dilute sulfuric acid. Addition of an oxidant such as sodium chlorate insured conversion of uranium minerals to a soluble form. At this point, the slurry was chemically reduced by a sulfide, such as sodium hydrosulfide, to remove substances such as ferric iron and molybdenum which were partially extracted by the solvent. The solids-liquid separation was accomplished by a four-stage counter cur rent decantation (CCD) thickener circuit. The uranium was extracted from pregnant liquor with 0.1M DDPA in kerosene and stripped with ION

Jan 1, 1961

By R. W. Heckel, R. L. DeGregorio

The DeHoff method of determining the size distribution of ellipsoid-shaped, second-phase particles has been applied to the spheroidization of cementite in a eutectoid steel. The surface area of Precipitates determined from the various size fractions in the distribution was correlated with surface-area measurements based on the number of intersected interfaces on a random lest line. The precipitate particles were found to be oblate ellipsoids with an axial ratio of about 0.90. The size distributions were found to be log-normal. A method is proposed whereby the shrinkage of small particles and growth of large ones can be determined from the experimental data. The experimental data are compared to previously proposed mathematical models describing diffusion-controlled kinetics and various types of interface-controlled kinetics. The experimental growth and shrinkage rates are considerably slower than those predicted by diffusion-controlled kinetics. The best fit is obtained for a model describing interface-controlled kinetics limited by the rate of formation of cementite at the growing interfaces where the interfacial reaction rate is proportional to the solute thermodynamic activity gradient across the surface. THE subject of spheroidization of second-phase particles has been considered previously by other investigators. Livingston1 has studied the precipitation of the cobalt-rich phase from copper alloys containing 0.7 to 3.2 wt pet Co. His results indicate that the average particle size increases as the cube root of the heat treatment time in accord with diffusion-controlled kinetics.&apos;-&apos; Komatsu and rant&apos; in their studies of the growth of SiO2 in a dispersion-strengthened copper-silica alloy found that the initial growth of SiO2 proceeded by diffusion-controlled growth and later stages were limited by interface-controlled growth. The values of the activation energy obtained for the interface- controlled process led them to the conclusion that the process was limited by the dissociation of SiO2 at the shrinking interface. The transition from diffusion-controlled growth to interface-controlled growth was characterized by a particle size which varied with the heat-treatment temperature. It is also interesting to note that the particle-size distribution as found by Komatsu and Grant exhibits log-normal behavior when replotted on log-probability coordinates. Dromsky, Lenel, and Ansell9 have observed the growth of Al2O3 particles having a mean free path of about 1.5 to 13 µ. Their photomicrographs indicate that the Al2O3 particle sizes were larger than those observed by Komatsu and Grant. Dromsky, Lenel, and Ansell concluded that the Al2O3 coarsened by an interface-controlled growth mechanism limited by the solution of A12O3 at the shrinking interfaces. Bannyh, Modin, and Modin10 have studied the spheroidization of a eutectoid steel (0.83 wt pet C). Their spheroidization (tempering) treatments were carried out in the range from 210o to 700°C for times between 1.5 sec and 20 hr. Measurements of mean particle size as a function of time indicated a cube root of time dependence of size at 700°C, in agreement with previous analyses of diffusion-controlled kinetics.2°7 At lower temperatures, the time dependence was less than the one-third power. It is important to note that, although mathematical models of growth* considering particle-size distribution have been available, measurements of only mean particle size have been carried out. DeHoff11 has presented a quantitative metallography technique which is applicable to the determination of the size distribution of ellipsoids of constant shape. This method is applicable to both oblate and prolate ellipsoids of all ratios of minor to major axis and is based upon an extension of Saltykov&apos;s analysis12 of the distribution of spheres of varying size. DeHoff&apos;s analysis is based upon measurements of the minor axis of the particles on a random plane of polish in a unit area. This method provides a measure of the number of ellipsoids in various size ranges per unit area. As pointed out by DeHoff, such measurements may be used to obtain surface area and total precipitate volume data. Thus, the accuracy of the distribution analysis may be checked by comparison of the surface and volume

Jan 1, 1965